High-strength steel plate and manufacturing method thereof

ABSTRACT

Disclosed is a high-strength sheet including: C: 0.15% by mass to 0.35% by mass, total of Si and Al: 0.5% by mass to 3.0% by mass, Mn: 1.0% by mass to 4.0% by mass, P: 0.05% by mass or less, and S: 0.01% by mass or less, with the balance being Fe and inevitable impurities, wherein the steel structure satisfies that: a ferrite fraction is 5% or less, the total fraction of tempered martensite and tempered bainite is 60% or more, the amount of retained austenite is 10% or more, MA has an average size of 1.0 μm or less, a half-width of the concentration distribution of Mn in the carbon-concentrated region that is equal to the amount of retained austenite is 0.3% by mass or more, and a scattering intensity at the q value of 1 nm −1  in X-ray small angle scattering is 1.0 cm −1  or less.

TECHNICAL FIELD

The present disclosure relates to a high-strength sheet that can be used in various applications including automobile parts.

BACKGROUND ART

In order to realize both weight reduction and collision safety, steel sheets used for automobile parts and the like are required to achieve both improvement in strength and improvement in impact resistance properties.

For example, Patent Document 1 discloses a high-strength steel sheet in which an attempt is made to improve impact resistance properties by heating a slab to 1,210° C. or higher and controlling the hot-rolling conditions to form fine TiN particles having a size of 0.5 μm or less, thereby suppressing the formation of AlN particles having a particle size of 1 μm or more that act as a starting point of low temperature fracture.

Patent Document 2 discloses a high-strength sheet in which an attempt is made to improve collision resistance properties by forming a network structure in which 50% or more of a ferrite grain size is in contact with a hard phase while defining the C amount to more than 0.45% and 0.77% or less, the Mn amount to 0.1% or more and 0.5% or less and the Si amount to 0.5% or less, and defining each addition amount of Cr, Al, N and O.

Patent Document 3 discloses a high-strength sheet in which an attempt is made to improve collision resistance properties by adding 3.5 to 10% of Mn, thereby adjusting the amount of retained austenite to 10% or more and an average interval of retained austenite to 1.5 μm or less.

Patent Document 4 discloses a high-strength sheet that has a tensile strength of 980 to 1,180 MPa and also exhibits satisfactory deep drawability.

PRIOR ART DOCUMENT Patent Document

Patent Document 1: JP 5240421 B1

Patent Document 2: JP 2015-105384 A

Patent Document 3: JP 2012-251239 A

Patent Document 4: JP 2009-203548 A

DISCLOSURE OF THE INVENTION Means or Solving the Problems

In order to realize further weight reduction, steel sheets used for automobile parts are required to have sufficient strength and impact resistance properties while being made thinner. Thus, steel sheets having higher tensile strength and excellent impact properties are required.

In various applications including automobile parts, steel sheets are required to have not only high tensile strength and impact properties, but also excellent strength-ductility balance, high yield ratio, stretch formability and excellent hole expansion ratio.

Specifically, the followings are required for each of the tensile strength, the strength-ductility balance, the yield ratio, the deep drawing properties and the hole expansion ratio.

The tensile strength is required to be 980 MPa or higher. In order to increase stress that can be applied during use, there is a need to have high yield strength (YS), in addition to high tensile strength (TS). From the viewpoint of ensuring collision safety or the like, there is a need to increase the yield strength of the steel sheet and to attain properties capable of suppressing fracture during deformation in order to stably exhibit strength properties upon collision. Therefore, specifically, there is required, together with a yield ratio (YR=YS/TS) of 0.75 or more, an increase in thickness reduction ratio of the fracture portion during a tensile test as an evaluation index substituting fracture properties. A joint strength of the spot welded portion is also required as basic performances of the steel sheet for automobiles. Specifically, a cross tensile strength of the spot welded portion is required to be 6 kN or higher.

Regarding the strength-ductility balance, the product (TS×EL) of TS and total elongation (EL) is required to be 20,000 MPa % or higher. In order to ensure the formability during parts forming, it is also required that the hole expansion ratio λ showing hole expansion properties is 20% or more and the maximum forming height (forming height) showing stretch formability is 16 mm or more.

However, it is difficult for the high-strength sheets disclosed in Patent Documents 1 to 4 to satisfy all of these requirements, and there has been required a high-strength steel sheet that can satisfy all of these requirements.

The embodiment of the present invention has been made to respond to these requirements, and it is an object thereof to provide a high-strength sheet in which all of tensile strength (TS), yield ratio (YR), the product (TS×EL) of (TS) and total elongation (EL), hole expansion ratio (λ), thickness reduction ratio (RA) of the fracture portion during a tensile test, maximum forming height and cross tensile strength (SW cross tension) of the spot welded portion are at a high level, and a manufacturing method thereof.

Means for Solving the Problems

Aspect 1 of the present invention provides a high-strength sheet including:

C: 0.15% by mass to 0.35% by mass,

total of Si and Al: 0.5% by mass to 3.0% by mass,

Mn: 1.0% by mass to 4.0% by mass,

P: 0.05% by mass or less, and

S: 0.01% by mass or less, with the balance being Fe and inevitable impurities,

in which the steel structure satisfies that:

a ferrite fraction is 5% or less,

the total fraction of tempered martensite and tempered bainite is 60% or more,

the amount of retained austenite is 10% or more,

MA has an average size of 1.0 μm or less,

a half-width of the concentration distribution of Mn in the carbon-concentrated region that is equal to the amount of retained austenite is 0.3% by mass or more, and

a scattering intensity at the q value of 1 nm⁻¹ in X-ray small angle scattering is 1.0 cm⁻¹ or less.

Aspect 2 of the present invention provides the high-strength sheet according to aspect 1, in which the C amount is 0.30% by mass or less.

Aspect 3 of the present invention provides the high-strength sheet according to aspect 1 or 2, in which the Al amount is less than 0.10% by mass.

Aspect 4 of the present invention provides a method for manufacturing a high-strength sheet, which includes:

preparing a rolled material including: C: 0.15% by mass to 0.35% by mass, total of Si and Al: 0.5% by mass to 3.0% by mass, Mn: 1.0% by mass to 4.0% by mass, P: 0.05% by mass or less, and S: 0.01% by mass or less, with the balance being Fe and inevitable impurities;

holding the rolled material at a temperature between an Ac₁ point and 0.2×Ac₁ point+0.8×Ac₃ point for 5 seconds or more, followed by heating to a temperature of an Ac₃ point or higher and further holding for 5 to 600 seconds, thereby austenitizing the rolled material,

after the austenitizing, cooling the material from a temperature of 650° C. or higher to a cooling stopping temperature between 100° C. or higher and lower than 300° C. at an average cooling rate of 10° C./sec or more;

heating the material from the cooling stopping temperature to a reheating temperature in a range of 300° C. to 500° C. at an average heating rate of 30° C./sec or more;

holding at the reheating temperature T so as to satisfy a tempering parameter P of 10,000 to 14,500 defined in the equation (1) and a holding time of 1 to 300 seconds; and

after the holding, cooling from the reheating temperature to 200° C. at an average cooling rate of 10° C./sec or more.

P=T×(20+log(t/3600))   (1)

where T: reheating temperature (K) and t: holding time (seconds).

Aspect 5 of the present invention provides the manufacturing method according to aspect 4, wherein cooling to the cooling stopping temperature includes:

cooling to a rapid cooling starting temperature that is a temperature of 650° C. or higher at an average cooling rate of 0.1° C./sec or more and less than 10° C./sec; and

cooling from the rapid cooling starting temperature to the cooling stopping at an average cooling rate of 10° C./sec or more.

Aspect 6 of the present invention provides the manufacturing method according to aspect 4 or 5, in which the tempering parameter is 11,000 to 14,000 and the holding time is 1 to 150 seconds.

Effects of the Invention

According to the embodiment of the present invention, it is possible to provide a high-strength sheet in which all of tensile strength (TS), yield ratio (YR), the product (TS×EL) of (TS) and total elongation (EL), hole expansion ratio (λ), thickness reduction ratio (RA) of the fracture portion during a tensile test (impact resistance properties), maximum forming height and cross tensile strength (SW cross tension) of the spot welded portion are at a high level, and a manufacturing method thereof.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1 is a diagram explaining a method for manufacturing a high-strength sheet according to the embodiment of the present invention, especially a heat treatment.

MODE FOR CARRYING OUT THE INVENTION

The inventors of the present application have intensively studied and found that it is possible to obtain a high-strength sheet in which all of tensile strength (TS), yield ratio (YR), the product (TS×EL) of (TS) and total elongation (EL), hole expansion ratio (λ), thickness reduction ratio (RA) of the fracture portion during a tensile test (impact resistance properties), maximum forming height and cross tensile strength (SW cross tension) of the spot welded portion are at a high level by allowing the steel structure (metal structure) to satisfy that: a ferrite fraction is 5% or less, the total fraction of tempered martensite and tempered bainite is 60% or more, the amount of retained austenite (γ) is 10% or more, MA has an average size of 1.0 μm or less, a half-width of the concentration distribution of Mn in the carbon-concentrated region that is equal to the amount of retained austenite is 0.3% by mass or more, and a scattering intensity at the q value of 1 nm⁻¹ in X-ray small angle scattering is 1.0 cm⁻¹ or less, in a steel including predetermined components.

As mentioned in detail later, the high-strength sheet according to the embodiment of the present invention has the Mn-concentrated region formed by holding in a two-phase coexistence region intermediate between an Ac₁ point and an Ac₃ point, more specifically at a temperature between the Ac₁ point and 0.2×Ac₁ point+0.8×Ac₃ point for a predetermined time, followed by holding at a temperature of the Ac₃ point or higher for a predetermined time, in the austenitizing step of a heat treatment during manufacturing. During the heat treatment, the carbon-concentrated region corresponding to retained austenite (the same amount as that of retained austenite) is formed. In this carbon-concentrated region, both the Mn-concentrated region and the Mn-nonconcentrated region are formed. In other words, the region containing a large amount of Mn and the region that does not contain a large amount of Mn exist in some carbon-enriched regions (retained austenite). Therefore, when the distribution of the Mn concentration is measured with respect to the entire carbon-concentrated region (i.e., corresponding to the entire retained austenite), there is a certain degree or more of variation in Mn concentration. Specifically, the half-width of the concentration distribution of Mn becomes 0.3% by mass or more.

Variation in amount of Mn contained in retained austenite means that it is possible to be provided with retained austenite with various degrees of stability. Therefore, residual austenite with low stability that causes strain induced transformation at a relatively small amount of strain and residual austenite with high stability that causes strain induced transformation at a large amount of strain coexist, thus making it possible to cause strain induced transformation in various strain regions. As a result, the n value can be increased in a wide strain region to enhance strain dispersibility, thus enabling realization of high stretch formability.

The high-strength sheet according to the embodiment of the present invention and a manufacturing method thereof will be described in detail below.

1. Steel Structure

The steel structure of the high-strength sheet according to the embodiment of the present invention will be described in detail below.

In the following description of the steel structure, there are cases where mechanisms capable of improving various properties by having such the structure are described. It should be noted that these mechanisms are those envisaged by the inventors of the present application based on the findings currently obtained, but do not limit the technical scope of the present invention.

(1) Ferrite Fraction: 5% or Less

Ferrite generally has excellent workability but has a problem such as low strength. As a result, a large amount of ferrite leads to a decrease in yield ratio. Therefore, a ferrite fraction was set at 5% or less (5 volume % or less).

The ferrite fraction is preferably 3% or less, and more preferably 1% or less.

The ferrite fraction can be determined by observing with an optical microscope and measuring the white region by the point counting method. By such a method, it is possible to determine the ferrite fraction by an area ratio (area %). The value obtained by the area ratio may be directly used as the value of the volume ratio (volume %).

(2) Total Fraction of Tempered Martensite and Tempered Bainite: 60% or More

By setting the total fraction of tempered martensite and tempered bainite at 60% or more (60 volume % or more), it is possible to achieve both high strength and high hole expansion properties. The total fraction of tempered martensite and tempered bainite is preferably 70% or more.

It is possible to determine the amounts of tempered martensite and tempered bainite (total fraction) by performing SEM observation of a Nital-etched cross-section, measuring a fraction of MA (i.e., the total of retained austenite and martensite as quenched) and subtracting the above-mentioned ferrite fraction and MA fraction from the entire steel structure.

(3) Amount of Retained Austenite: 10% or More

The retained austenite causes the TRIP phenomenon of being transformed into martensite due to strain induced transformation during working such as press working, thus making it possible to obtain large elongation. Martensite thus formed has high hardness. Therefore, excellent strength-ductility balance can be obtained. By setting the amount of retained austenite at 10% or more (10 volume % or more), it is possible to realize TS×EL of 20,000 MPa % or more and excellent strength-ductility balance.

The amount of retained austenite is preferably 15% or more.

In the high-strength sheet according to the embodiment of the present invention, most of retained austenite exists in the form of MA. MA is abbreviation of a martensite-austenite constituent and is a composite (complex structure) of martensite and austenite.

It is possible to determine the amount of retained austenite by calculating a diffraction intensity ratio of ferrite (including tempered martensite and untempered martensite in X-ray diffraction) and austenite by X-ray diffraction, followed by calculation. As an X-ray source, Co—Kα ray can be used.

(4) Average Size of MA: 1.0 μm or Less

MA is a hard phase and the vicinity of matrix/hard phase interface acts as a void forming site during deformation. The larger the MA size, the more strain concentration occurs at the matrix/hard phase interface, thus easily causing fracture from voids formed in the vicinity of the matrix/hard phase interface as a starting point.

Therefore, it is possible to improve the hole expansion ratio λ by decreasing the MA size, especially the MA average size to 1.0 μm or less, thereby suppressing fracture.

The average size of MA is preferably 0.8 μm or less.

It is possible to determine the average size of MA by observing a Nital-etched cross-section in three or more fields of view at a magnification of 3,000 times or more with SEM, drawing a straight line of 200 μm or more in arbitrary position in the micrograph, measuring the length of intercept where the straight line crosses MA, and calculating the average of the intercept lengths.

(5) Half-Width of Concentration Distribution of Mn in Carbon-Concentrated Region that is Equal to Amount of Retained Austenite is 0.3% by Mass or More

As mentioned above, most of retained austenite exists in the form of MA, and it is difficult to identify only by retained austenite using a microscope or SEM. In retained austenite, since the solid solubility limit of carbon is larger than that of ferrite or the like, carbon is concentrated in retained austenite by performing a heat treatment mentioned later. Therefore, element mapping of carbon is performed using an electron probe micro analyzer (EPMA) and the measurement point having the amount equal to that of retained austenite determined by the above-mentioned X-ray diffraction in order of the measurement point having ascending carbon concentration was taken as the carbon-concentrated region, and thus this carbon-concentrated region can be judged as retained austenite. That is, for example, when the amount of retained austenite is 15% by volume, by selecting 15% of measurement points having higher carbon concentration from the measuring points whose carbon content was measured by element mapping, it is possible to judge that each of these measurement points having high carbon concentration (carbon-concentrated region) is retained austenite.

Therefore, “carbon-concentrated region that is equal to the amount of retained austenite” means the region that corresponds to retained austenite.

It is also possible to measure the concentration distribution of Mn in the carbon-concentrated region that is equal to the amount of retained austenite, especially the half-width of the concentration distribution of Mn, using EPMA. The distribution of the Mn amount at the measurement point, that was judged as the carbon-concentrated region, is graphed, thus making it possible to obtain a half-width therefrom.

The larger the half-width of Mn of the concentration distribution, the larger the variation in the concentration of Mn in retained austenite (the range of the concentration distribution of Mn is wide) becomes. In the high-strength sheet according to the embodiment of the present invention, the half-width of the concentration distribution of Mn is 0.3% by mass or more, preferably 0.5% by mass or more, more preferably 0.6% by mass or more, and still more preferably 0.75% by mass or more.

In this way, by varying the amount of Mn contained in retained austenite (carbon-enriched region), it is possible to form retained austenite with a wide range of stability from retained austenite with low stability to retained austenite with high stability. Retained austenite with low stability causes strain induced transformation at a small amount of strain, thus turning to martensite. Retained austenite with high stability does not cause strain induced transformation at a small amount of strain, and causes strain induced transformation, thus turning martensite only after a large amount of strain is applied. Therefore, when retained austenite having a wide range of stability exists, retained austenite continuously causes strain induced transformation from the time when a small amount of strain is applied immediately after the start of forming to the time when forming progresses and a large amount of strain is applied. As a result, the n value can be increased over a wide range of strain to enhance the strain dispersibility, thus enabling realization of high stretch formability.

(6) Scattering Intensity at q Value of 1 nm⁻¹ in X-Ray Small Angle Scattering: 1.0 cm⁻¹ or Less

X-ray small angle scattering means that the size distribution of fine particles (e.g., cementite particles dispersed in a steel sheet) contained in the steel sheet can be obtained by irradiating the steel sheet with X-rays and measuring scattering of X-rays transmitted through the steel sheet. In the steel sheet according to the embodiment of the present invention, it is possible to determine the size distribution of cementite particles that are fine particles dispersed in tempered martensite by X-ray small angle scattering. Specifically, in X-ray small angle scattering, it is possible to analyze the size and the fraction of cementite particles using the q value and the scattering intensity.

The q value is an index of the size of particles (e.g., cementite particles) in the steel sheet. The “q value of 1 nm⁻¹” corresponds to cementite particles having a particle size of about 1 nm. The scattering intensity is an index of the volume fraction of particles (e.g., cementite particles) in the steel sheet. The larger the scattering intensity, the larger the volume fraction of cementite becomes.

The scattering intensity at a certain q value semi-quantitatively indicates the volume fraction of cementite particles of the size corresponding to the q value. For example, the scattering intensity at the q value of 1 nm⁻¹ semi-quantitatively indicates the volume fraction of fine cementite particles having a size of about 1 nm.

In other words, large scattering intensity at the q value of 1 nm⁻¹ indicates large volume fraction of fine cementite particles having a size of about 1 nm. In the steel sheet in which “the scattering intensity at the q value of 1 nm⁻¹ is 1.0 cm⁻¹ or less”, it means that the volume fraction of fine cementite particles having a size of about 1 nm existing in the steel sheet is a predetermined value (the value corresponding to the scattering intensity of 1.0 cm⁻¹) or less. As described later, it is considered that the steel sheet in which “the scattering intensity at the q value of 1 nm⁻¹ is 1.0 cm⁻¹ or less” is excellent in collision resistance properties since the volume fraction of cementite having a size of about 1 nm is suppressed to a low value.

In high-ductility steel containing retained γ, it is preferable that no cementite ideally exists in a state where carbon is concentrated in retained austenite. Fine cementite having a grain size of about 1 nm dispersed in the steel material hinders dislocation migration, thus enabling degradation of the deformability of the steel material. Therefore, in the steel material having a large volume fraction of cementite having a grain size of about 1 nm, fracture during deformation is promoted, thus enabling degradation of collision resistance properties.

In the steel sheet according to the embodiment of the present invention, by suppressing the volume fraction of fine cementite to a low value, more specifically, by setting the scattering intensity at the q value of 1 nm⁻¹ at 1 cm⁻¹ or less, fine carbide formed in laths of tempered martensite is reduced to enhance the deformability in martensite. Thus, fracture of the steel sheet upon collision is suppressed to improve collision resistance properties of the steel sheet.

X-ray small angle scattering was measured using a Nano-viewer, Mo tube manufactured by Rigaku Corporation. A 3 mmφ disk-shaped sample was cut out from the steel sheet and samples having a thickness of 20 μm were cut out from the vicinity of the thickness of ¼ and then used. Data at the q value of 0.1 to 10 nm⁻¹ were collected. Among them, absolute intensity was determined for the q value of 1 nm⁻¹.

(7) Other Steel Structure:

In the present description, steel structures other than the above-mentioned ferrite, tempered martensite, tempered bainite retained austenite and cementite are not specifically defined. However, pearlite, untempered bainite, untempered martensite and the like may exist, in addition to the steel structures such as ferrite. As long as the steel structure such as ferrite satisfies the above-mentioned structure conditions, the effects of the present invention are exhibited even if pearlite or the like exists in the steel.

2. Composition

The composition of the high-strength sheet according to the embodiment of the present invention will be described below. Main elements C, Si, Al, Mn, P and S will be described. Note that all percentages as unit with respect to the composition are by mass.

(1) C: 0.15 to 0.35%

Carbon (C) is an element indispensable for ensuring properties such as high strength-ductility balance (TS×EL balance) by increasing the amount of desired structure, especially retained y. In order to effectively exhibit such effect, there is a need to add C in the amount of 0.15% or more. However, the amount of more than 0.35% is not suitable for welding. The amount is preferably 0.18% or more, and more preferably 0.20% or more. The amount is preferably 0.30% or less. If the C amount is 0.25% or less, welding can be easily performed.

(2) Total of Si and Al: 0.5 to 3.0%

Si and Al each have the effect of suppressing the precipitation of cementite, thus remaining retained austenite. In order to effectively exhibit such effect, there is a need to add Si and Al in the total amount of 0.5% or more. If the total amount of Si and Al exceeds 3.0%, the deformability of the steel is degraded, thus degrading TS×EL and forming height. The total amount is preferably 0.7% or more, and more preferably 1.0% or more. The total amount is preferably 2.5% or less.

Note that Al may be added in the amount enough to function as an deoxidizing element, i.e., less than 0.10% by mass. For the purpose of suppressing the formation of cementite to increase the amount of retained austenite, Al may be added in a larger amount of 0.7% by mass or more.

(3) Mn: 1.0 to 4.0%

Mn suppresses the formation of ferrite. Mn is an element indispensable for improving the stretch formability by forming the Mn-concentrated region and forming retained austenites with different stabilities. In order to effectively exhibit such effect, there is a need to add Mn in the amount of 1.0% or more. If the amount exceeds 4.0%, it is difficult to control because of narrow temperature range of two-phase region heating, and transformation does not proceed even by holding at a temperature between the Ac₁ point and 0.2×Ac₁ point+0.8×Ac₃ point for a predetermined time holding because of too lowered temperature, thus failing to form the Mn-concentrated region in some cases. The amount is preferably 1.5% or more, and more preferably 2.0% or more. The amount is preferably 3.5% or less.

(4) P: 0.05% or Less

P inevitably exists as an impurity element. If more than 0.05% of P exists, EL and λ are degraded. Therefore, the content of P is set at 0.05% or less (including 0%). Preferably, the content is 0.03% or less (including 0%).

(5) S: 0.01% or Less

S inevitably exists as an impurity element. If more than 0.01% of S exists, sulfide-based inclusions such as MnS are formed and act as a starting point of cracking, thus degrading λ. Therefore, the content of S is set at 0.01% or less (including 0%). The content is preferably 0.005% or less (including 0%).

(6) Balance

In a preferred embodiment, the balance is composed of iron and inevitable impurities. It is permitted to mix, as inevitable impurities, trace elements (e.g., As, Sb, Sn, etc.) incorporated according to the conditions of raw materials, materials, manufacturing facilities and the like. There are elements whose content is preferably as small as possible, like P and S, that are therefore inevitable impurities in which the composition range is separately defined as mentioned above. Therefore, “inevitable impurities” constituting the balance as used herein means the concept excluding elements whose composition range is separately defined.

However, it is not limited to this embodiment. As long as properties of the high-strength steel sheet according to the embodiment of the present invention can be maintained, any other element may be further included.

3. Properties

As mentioned above, regarding the high-strength sheet according to the embodiment of the present invention, all of TS, YR, TS×EL, λ, collision resistance properties and SW cross tension are at a high level. These properties of the high-strength sheet according to the embodiment of the present invention will be described in detail below.

(1) Tensile Strength (TS)

The high-strength sheet has TS of 980 MPa or higher. Preferably, TS is 1,180 MPa or higher. If TS is lower than 980 MPa, excellent fracture properties can be more surely obtained, but it is not preferable since withstand load upon collision decreases.

(2) Yield Ratio (YR)

The high-strength sheet has an yield ratio of 0.75 or more. This makes it possible to realize a high yield strength combined with the above-mentioned high tensile strength and to use the final product obtained by working such as deep drawing under high stress. Preferably, the high-strength sheet has a yield ratio of 0.80 or more.

(3) The Product (TS×EL) of TS and Total Elongation (EL)

TS×EL is 20,000 MPa % or more. By having TS×EL of 20,000 MPa % or more, it is possible to obtain high-level strength-ductility balance that has both high strength and high ductility. Preferably, TS×EL is 23,000 MPa % or more.

(4) Stretch Formability (Maximum Forming Height)

The maximum forming height is an index used for evaluation of the stretch formability. The maximum forming height is taken as a punch stroke at the occurrence of fracture where the load rapidly decreases in the load-stroke diagram.

More specifically, using a 120 mmφ test piece and using a 53.6 mmφ die having a shoulder radius of 8 mm and a 50 mmφ spherical head punch, a lubricating polyethylene sheet is interposed between the punch and the steel sheet, and stretch forming is performed at a blank holding force of 1,000 kgf, and then the maximum forming height is determined by measuring the height at the time of fracture (punch stroke).

The high-strength steel sheet according to the embodiment of the present invention has the maximum forming height of 20 mm or more, and preferably 21 mm or more.

(5) Hole Expansion Ratio (λ)

A hole expansion ratio λ is determined in accordance with Japan Iron and Steel Federation Standard JFS T1001. A punched hole having a diameter d₀ (d₀=10 mm) was formed in a test piece and a punch having a tip angle of 60° was pushed into this punched hole, and a diameter d of the punched hole at the time when the generated cracking penetrated the thickness of the test piece was measured, and then the hole expansion ratio is calculated by the following equation.

λ(%)={(d−d ₀)/d ₀}×100

The high-strength sheet according to the embodiment of the present invention has a hole expansion ratio λ of 20% or more, and preferably 30% or more. This makes it possible to obtain excellent workability such as press formability.

(6) Thickness Reduction Ratio in Tensile Test (R5 Tensile Thickness Reduction Ratio)

Using a test piece provided with an arcuate notch having a radius of 5 mm on a No. 5 test piece, a tensile test was performed at a deformation rate of 10 mm/min and the test piece was fractured. Then, the fracture surface was observed and the value (t₁/t₀) obtained by dividing a thickness t₁ in a thickness direction of the fracture surface by an original thickness t₀ was taken as a thickness reduction ratio.

The thickness reduction rate in this test is 50% or more, preferably 52% or more, and more preferably 55% or more. This makes it possible to obtain a steel sheet having excellent impact resistance properties since the steel sheet is hardly fractured even if it deforms greatly upon collision.

(7) Cross Tensile Strength of Spot Welding

Cross tensile strength of spot welding was evaluated in accordance with JIS Z 3137. Two 1.4 mm-thick steel sheets laid one upon another were used as the conditions of spot welding. Using a dome radius type electrode, spot welding was performed under a welding pressure of 4 kN by increasing a current by 0.5 kA in a range from 6 kA to 12 kA, and the current value (minimum current value) at which dust was generated during welding was examined. A cross joint spot-welded at a current that is 0.5 kA lower than the minimum current value was used as a test piece for measurement of a cross tensile strength. Samples having a cross tensile strength of 6 kN or higher were rated “Good”. The cross tensile strength is preferably 8 kN or higher, and more preferably 10 kN or higher.

When the cross tensile strength is 6 kN or higher, it is possible to obtain parts having high bonding strength during welding when automobile parts and the like are manufactured from the steel sheet.

4. Manufacturing Method

The method for manufacturing a high-strength sheet according to the embodiment of the present invention will be described below.

The inventors of the present application have found that the above-mentioned desired steel structure is attained by subjecting a rolled material with predetermined composition to a heat treatment (multi-step austempering treatment) mentioned later, thus obtaining a high-strength steel sheet having the above-mentioned desired properties.

Details will be described below.

FIG. 1 is a diagram explaining a method for manufacturing a high-strength sheet according to the embodiment of the present invention, especially a heat treatment.

The rolled material to be subjected to the heat treatment is usually produced by cold-rolling after subjecting to hot-rolling. However, the process is not limited thereto, and the rolled material may be produced by any one of hot-rolling and cold-rolling. The conditions of hot-rolling and cold-rolling are not particularly limited.

(1) Austenitizing

In the austenitizing step, as shown in [2] of FIG. 1, a rolled material is held in a two-phase coexistence region intermediate between an Ac₁ point and an Ac₃ point, more specifically at a temperature T₁ (Ac₁≤T₁≤0.2×Ac₁ point+0.8×Ac₃) between the Ac₁ point and 0.2×Ac₁ point+0.8×Ac₃ point for 5 seconds or more and, as shown in [3] and [4] of FIG. 1, the rolled material is held at a heating temperature T₂ to a temperature T₂ (Ac₃≤T₂) of the Ac₃ point or higher for 5 to 600 second, thus austenitizing the rolled material.

The rolled material is heated to the temperature T₁, followed by holding for 5 seconds or more. The holding time is preferably 900 seconds or less. The holding temperature T₁ may be held at a constant temperature between the Ac₁ point and 0.2×Ac₁ point+0.8×Ac₃ point, as shown in [2] of FIG. 1, and the holding temperature may be varied between the Ac₁ point and 0.2×Ac₁ point+0.8×Ac₃ point, for example, slow heating is performed at a temperature between the Ac₁ point and 0.2×Ac₁ point+0.8×Ac₃ point. In this way, by holding in a relatively low temperature range within the two-phase coexistence region of ferrite and austenite, a large amount of Mn is distributed to the austenite side among coexisting ferrite and austenite, thus making it possible to obtain the Mn-concentrated region. Since austenite formed in the Mn-concentrated region remaining as retained austenite after the heat treatment has high Mn concentration, it is possible to increase the variation in the concentration of Mn in the carbon-concentrated region, thus enabling realization of high stretch formability.

If the temperature T₁ is lower than the Ac₁ point, the amount of Mn-concentrated austenite decreases to reduce the variation in the concentration of Mn in retained austenite (carbon-concentrated region), thus failing to obtain sufficient stretch formability.

If the temperature T₁ is higher than 0.2×Ac₁ point+0.8×Ac₃ point, the Mn concentration of austenite decreases to reduce the variation in the concentration of Mn in retained austenite (carbon-concentrated region), thus failing to obtain sufficient stretch formability.

If the holding time at the temperature T₁ is less than 5 seconds, because of insufficient Mn diffusion time, the concentration of Mn into austenite becomes insufficient to reduce the variation of Mn in retained austenite (carbon concentrated region), thus failing to obtain sufficient overhang formability.

The longer the holding time at the temperature T1, the better, and the holding time is preferably 900 seconds or less from the viewpoint of the productivity.

The temperature T₁ is preferably between 0.9×Ac₁ point+0.1×Ac₃ point and 0.3×Ac₁ point+0.7×Ac₃ point, and the holding time at the temperature T₁ is preferably 10 seconds or more and 800 seconds or less. The temperature T₁ is more preferably between 0.8×Ac₁ point+0.2×Ac₃ point and 0.4×Ac₁ point+0.6×Ac₃ point, and the holding time at the temperature T₁ is more preferably 30 seconds or more and 600 seconds or less.

The heating rate to the temperature T₁ shown as [1] in FIG. 1 is preferably 5 to 20° C./sec.

Next, as shown in [3] and [4] of FIG. 1, a material is heated to a temperature T₂ (Ac₃≤T₂) of an Ac₃ point or higher, followed by holding at the temperature T₂, thus austenitizing the rolled material. The holding time at the temperature T₂ is 5 to 600 seconds.

By heating to the temperature T₂ of the Ac₃ point or higher, the portion which was ferrite turns to austenite when heated to the temperature T₁. In this portion newly transformed into austenite, Mn is not concentrated. Therefore, the Mn-nonconcentrated region exists in austenite, together with the Mn-concentrated region, and it becomes possible to increase the variation in the concentration of Mn in retained austenite (carbon-concentrated region) in the high-strength sheet after the heat treatment, thus enabling realization of the high stretch formability.

If the temperature T₂ is lower than the Ac₃ point or the holding time at the temperature T₂ is less than 5 seconds, the ferrite fraction of the obtained high-strength sheet exceeds 5%, leading to a decrease in YR.

If the temperature T₂ is too high, Mn in the previously formed Mn-enriched region may diffuse and the variation in the Mn concentration may become too small. Therefore, the temperature T₂ is preferably the Ac₃ point+50° C. or lower.

If the holding time at the temperature T₂ is more than 600 seconds, the Mn concentration in the Mn-concentrated region decreases to reduce the variation in the Mn concentration in retained austenite, thus degrading the stretch formability.

The temperature T₂ is preferably the Ac₃ point+10° C. or higher, and the holding time at the temperature T₂ is preferably 10 to 450 seconds. The temperature T₂ is more preferably the Ac₃ point+20° C. or higher, and the holding time at the temperature T₂ is more preferably 20 to 300 seconds.

Heating from the temperature T₁ to the temperature T₂ shown in [3] of FIG. 1 is preferably performed at a heating rate of 0.1° C./sec or more and less than 10° C./sec.

The Ac₁ point and the Ac₃ point may be determined by the measurement, or may be calculated by a generally known calculation formula using the composition.

For example, the Ac₁ point and the Ac₃ point can be calculated using the following equations (1) and (2) (see, for example, Leslie, “The Physical Metallurgy of Steels”, Maruzen Co., Ltd., 1985):

Ac ₁ point (° C.)=723+29.1×[Si]−10.7×[Mn]+16.9×[Cr]−16.9×[Ni]  (1)

Ac ₃ point (° C.)=910−203×[C]^(1/2)+44.7×[Si]−30×[Mn]+700×[P]+400×[Al]+400×[Ti]+104×[V]−11×[Cr]+31.5×[Mo]−20×[Cu]−15.2×[Ni]  (2)

where the parentheses represent the content expressed by % by mass therein. (2) Cooling to Cooling Stopping Temperature Between 100° C. or Higher and Lower than 300° C.

After the austenitizing, as shown in [6] of FIG. 1, cooling is performed to a cooling stopping temperature T₃ between 100° C. or higher and lower than 300° C. at an average cooling rate of 10° C./sec or more. By controlling the cooling stopping temperature to a temperature in a range of 100° C. or higher and lower than 300° C., the final amount of retained austenite is controlled by adjusting the amount of austenite remaining without being transformed into martensite.

Cooling is performed at an average cooling rate of 10° C./sec or more between at least 650° C. and 300° C. This is because, by setting the average cooling rate at 10° C./sec or more, the formation of ferrite during cooling is suppressed to form a structure composed mainly of fine martensite.

Preferred example of such cooling includes cooling (slow cooling) to a rapid cooling starting temperature T₄ of 650° C. or higher at relatively low average cooling rate of 0.1° C./sec or more and 10° C./sec or less, as shown in [5] of FIG. 1, followed by cooling (rapid cooling) from the rapid cooling starting temperature T₄ to a cooling stopping temperature T₃ of 300° C. or lower at an average cooling rate of 10° C./sec or more and less than 200° C./sec, as shown in [6] of FIG. 1. By setting the rapid cooling starting temperature T₄ at 650° C. or higher, it is possible to suppress the formation of ferrite during cooling (slow cooling).

If the cooling rate is less than 10° C./sec, ferrite is formed to decrease YR. MA becomes coarse, thus decreasing the hole spreading ratio.

If the cooling stopping temperature T₃ is lower than 100° C., the amount of retained austenite becomes insufficient. As a result, TS increases but EL decreases, leading to insufficient TS×EL balance.

If the cooling stopping temperature T₃ is 300° C. or higher, coarse untransformed austenite increases and remains even after the subsequent cooling. Finally, the size of MA becomes coarse, thus decreasing the hole expansion ratio λ.

The cooling rate is preferably 15° C./° C. or higher, and more preferably 20° C./sec or more. The cooling stopping temperature T₃ is preferably 120° C. or higher and 280° C. or lower, and more preferably 140° C. or higher and 260° C. or lower.

As shown in [7] of FIG. 1, holding may be performed at the cooling stopping temperature T₃. In the case of holding, the holding time is preferably 1 to 150 seconds. Even if the holding time is more than 150 seconds, the productivity of the steel sheet is degraded though properties of the obtained steel sheet are not significantly improved. Therefore, the holding time is preferably set at 150 seconds or less.

(3) Reheating to Temperature in Range of 300° C. to 500° C.

As shown in [8] of FIG. 1, heating is performed from the above cooling stopping temperature to a reheating temperature in a range of 300° C. to 500° C. at a reheating rate of 30° C./sec or more. Rapid heating enables a decrease in retention time in a temperature range where precipitation and growth of carbide are promoted, thus making it possible to suppress the formation of fine carbide. The reheating rate is preferably 60° C./sec or more, and more preferably 70° C./sec.

Such rapid heating can be achieved by a method such as high-frequency heating or electric heating.

After reaching the reheating temperature T₅, as shown in [9] of FIG. 1, holding is performed at the same temperature. At that time, it is preferred that a tempering parameter P represented by the following equation (1) is set at 10,000 or more and 14,500 or less and the holding time is set at 1 to 150 seconds. The tempering parameter P of the steel sheet of the present embodiment is represented by the following equation (1):

P=T(K)×(20+log(t/3600)   (1)

where T is a tempering temperature (K) and t is a holding time (seconds).

During reheating, redistribution of carbon that is supersaturatedly sold-soluted in martensite occurs. Specifically, two phenomena, i.e. carbon diffusion from martensite to austenite and precipitation of carbide (cementite) in martensite laths. Among two phenomena, the precipitation of carbide easily occurs when holding is performed at low temperature for a long time. Even in the case of holding at high temperature, carbide is precipitated when the heating rate is low or the holding time is too long. Meanwhile, since carbon diffusion from martensite to austenite strongly depends on the diffusion rate, carbon diffusion can be sufficiently performed by a heat treatment at high temperature in a short time.

Particles of cementite existing in martensite easily act as a starting point of collision fracture and can degrade collision resistance properties. Therefore, in the case of reheating, it is desired that a reheating treatment is performed to promote carbon diffusion from martensite to austenite while suppressing the precipitation of carbide (cementite) in martensite laths. Thus, it is effective to perform rapid heating and a heat treatment at high temperature in a short time.

In order to obtain desired tensile strength by causing sufficient carbon diffusion, there is a need to control the tempering parameter P as a factor of a combination of temperature and time within a given range.

When the tempering parameter P is less than 10,000, carbon diffusion from martensite to austenite does not sufficiently occur and austenite becomes unstable, thus failing to ensure the amount of retained austenite, leading to insufficient TS×EL balance. If the tempering parameter P is more than 14,500, the formation of carbide cannot be prevented even by a short-time treatment, thus failing to ensure the amount of retained austenite, leading to degradation of TS×EL balance. Even if the tempering parameter is appropriate, carbide is formed in martensitic laths if the heating rate is too low and heating time is too long, so that crack propagation easily occurs during collision deformation, thus degrading collision resistance properties. The amount of carbide in martensite laths can be determined from the scattering intensity of X-ray small angle scattering.

If the reheating temperature T₅ is lower than 300° C., diffusion of carbon becomes insufficient, thus failing to obtain sufficient amount of retained austenite, leading to degradation of TS×EL. If the reheating temperature T₅ is higher than 500° C., retained austenite is decomposed into cementite and ferrite, thus failing to ensure properties because of insufficient retained austenite.

If holding is not performed or the holding time is less than 1 second, carbon diffusion may be insufficient, similarly. Therefore, it is preferred to hold at a reheating temperature T₅ for 1 second or more. If the holding time is more than 150 seconds, carbon may precipitate as cementite, similarly. Therefore, the holding time is preferably 150 seconds or less.

The reheating temperature T₅ is preferably 320 to 480° C., and more preferably 340 to 460° C.

The tempering parameter P is preferably 10,500 to 14,500, and the holding time at this time is preferably 1 to 150 seconds. The tempering parameter P is more preferably 11,000 to 14,000, and the holding time at this time is preferably 1 to 100 seconds, and more preferably 1 to 60 seconds.

After reheating, as shown in [10] of FIG. 1, cooling may be performed to the temperature of 200° C. or lower, for example, room temperature. The average cooling rate to 200° C. or lower is preferably 10° C./sec.

The high-strength sheet according to the embodiment of the present invention can be obtained by the above-mentioned heat treatment.

There is a possibility for person with ordinary skill in the art who came into contact with the method of manufacturing a high-strength steel sheet according to the embodiment of the present invention described above to obtain the high strength steel sheet according to the embodiment of the present invention by trial and error, using a manufacturing method different from the above-mentioned method.

EXAMPLES 1. Fabrication of Samples

After producing a cast material with the chemical composition shown in Table 1 was produced by vacuum melting, this cast material was hot-forged to form a steel sheet having a thickness of 30 mm and then hot-rolled. In Table 1, an Ac₃ point calculated from the composition was also shown.

Although the conditions of hot-rolling do not have a substantial influence on the final structure and properties of the present patent, a steel sheet having a thickness of 2.5 mm was produced by multistage rolling after heating to 1,200° C. At this time, the end temperature of hot-rolling was set at 880° C. After cooling to 600° C. at 30° C./sec, cooling was stopped and the steel sheet was inserted into a furnace heated to 600° C., held for 30 minutes and then furnace-cooled to obtain a hot-rolled steel sheet.

The hot-rolled steel sheet was subjected to pickling to remove the scale on the surface, and then cold-rolled to reduce the thickness to 1.4 mm. This cold rolled sheet was subjected to a heat treatment to obtain samples. The heat treatment conditions are shown in Table 2. The number in parentheses, for example, [2] in Table 2 corresponds to the process of the same number in parentheses in FIG. 1.

In Table 2, sample No. 1 is sample in which austenitizing was performed only at a temperature of the Ac₃ point or higher corresponding to a temperature T₂ without performing in two stages of austenitizing at a temperature T₁ and austenitizing at the temperature T₂.

Sample No. 9 is sample in which cooling was performed to a reheating temperature, followed by holding the same temperature instead of cooling to a cooling stopping temperature between 100° C. or higher and lower than 300° C. (samples in which the steps corresponding to [7] to [8] in FIG. 1 were skipped).

Samples 15 and 31 to 36 are samples in which the heating temperature T₂ and the rapid cooling starting temperature T₄ were set at the same temperature. After the austenitizing, cooling was performed to a cooling stopping temperature T₃ in a single stage.

Reheating corresponding to [8] was performed by an electric heating method.

In Table 1 to Table 4, the numerical value with an asterisk (*) indicates that it deviates from the range of the embodiment of the present invention.

TABLE 1 Composition C Si Mn P S Al Si + Al 0.2 Ac₁ + Steel % by % by % by % by % by % by % by Ac₁ AC₃ 0.8 Ac₃ No. mass mass mass mass mass mass mass ° C. ° C. ° C. a 0.32 1.61 1.95 0.011 0.001 0.04 1.65 749 826 810 b 0.22 2.10 1.80 0.005 0.001 0.03 2.13 765 868 847 c *0.10 1.42 2.51 0.008 0.002 0.03 1.45 737 847 825 d 0.19 1.28 *5.20 0.007 0.002 0.03 1.31 705 736 730 e 0.21 1.54 *0.63 0.010 0.002 0.03 1.57 761 880 856 f 0.27 0.20 2.19 0.011 0.002 0.03 *0.23 705 761 750 g *0.48 1.51 1.67 0.014 0.002 0.04 1.55 749 804 793 h 0.28 3.20 2.47 0.014 0.002 0.04 *3.24 790 888 868 i 0.26 0.91 2.35 0.007 0.002 0.03 0.94 724 790 777 j 0.28 1.52 1.70 0.007 0.002 0.04 1.56 749 837 819 k 0.22 1.34 1.50 0.013 0.001 0.04 1.38 746 846 826 l 0.21 1.70 2.03 0.014 0.003 0.02 1.72 751 841 823 m 0.20 1.32 2.25 0.013 0.001 0.03 1.35 737 823 806 n 0.26 1.19 1.52 0.013 0.002 0.03 1.22 741 828 811 o 0.28 0.84 1.88 0.007 0.002 0.25 1.09 727 885 853 p 0.26 0.91 1.90 0.011 0.002 0.02 0.93 729 799 785 q 0.26 1.36 2.02 0.007 0.003 0.04 1.40 741 824 808

TABLE 2 Heat treatment conditions [6] Rapid [6] [2] [4] cool- Cool- [8] Hold- Heat- [5] ing ing Reheat- ing [2] [3] ing [4] Slow start- [6] stopp- [7] [8] ing [9] [10] temper- Hold- Heat- temper- Hold- cool- ing Cool- ing Hold- Reheat- temper- Hold- Cool- ature ing ing ature ing ing temper- ing temper- ing ing ature ing ing Steel T, time rate T₂ time rate ature T₄ rate ature T₃ time rate T₅ time rate Param- No. No. ° C. Sec ° C./sec ° C. Sec ° C./sec ° C. ° C./sec ° C. Sec ° C./sec ° C. Sec ° C./sec eter 1 a *— *— — 850 120 10 700 28 200 50 100 400 *300 10 12734 2 a *700 100 20 850 20 10 700 28 200 50 100 440 20 10 12652 3 a *830 100 20 850 20 10 700 28 200 50 100 400 *300 10 12734 4 a 780 100 0 *780 20 10 700 28 200 50 100 440 20 10 12652 5 a 780 100 0 *780 20 10 700 28 275 50 100 440 20 10 12652 6 a 780 100 20 *800 20 10 700 28 125 50 100 440 20 10 12652 7 a 780 100 20 850 20 10 700 28 *350 50 100 440 20 10 12652 8 a 780 *3 20 850 20 10 700 28 150 50 100 440 20 10 12652 9 a 780 100 20 850 *950 10 700 28 *400 — *— 400 *300 10 12734 10 a 780 100 20 850 20 10 700 28 *20 50 100 440 20 10 12652 11 a 780 100 20 850 20 10 700 28 200 50 30 440 20 10 12652 12 a 780 100 20 850 20 10 700 28 200 50 90 400 50 10 12210 13 a 780 100 20 850 20 10 700 28 200 50 100 350 50 10 11303 14 a 780 100 20 850 20 10 700 28 200 50 100 350 20 10 11055 15 a 780 100 20 850 20 — 850 28 200 50 100 440 20 10 12652 16 a 780 100 20 850 20 10 *580 28 200 50 100 400 *300 10 12734 17 a 780 100 20 850 20 10 700 28 200 50 *15 400 20 10 11942 18 a 780 100 20 850 20 10 700 10 200 50 100 440 20 10 12652 19 a 780 100 20 850 20 10 700 28 200 50 100 *550 20 10 *14604 20 a 780 100 20 850 20 10 700 28 200 50 100 *250 20 10 *9280 21 b 810 100 20 900 20 10 700 28 200 50 100 400 20 10 11942 22 c 780 100 20 900 20 10 700 28 200 50 100 400 20 10 11942 23 d 720 100 20 850 20 10 700 28 200 50 100 400 20 10 11942 24 e 800 100 20 900 20 10 700 28 200 50 100 400 20 10 11942 25 f 740 100 20 850 20 10 700 28 200 50 100 400 20 10 11942 26 g 780 100 20 850 20 10 700 28 200 50 100 400 20 10 11942 27 h 840 100 20 940 20 10 700 28 200 50 100 400 20 10 11942 28 i *780 100 20 850 20 10 700 28 200 50 100 400 20 10 11942 29 j 780 100 20 850 20 10 700 28 200 50 100 400 20 10 11942 30 k 780 100 20 850 20 10 700 28 200 50 100 400 20 10 11942 31 l 780 100 20 850 20 — 850 28 200 50 100 400 20 10 11942 32 m 780 100 20 850 20 — 850 28 200 50 100 400 20 10 11942 33 n 780 100 20 850 20 — 850 28 200 50 100 400 20 10 11942 34 o 780 100 20 920 20 — 920 28 200 50 100 400 20 10 11942 35 p 780 100 20 850 20 — 850 28 200 50 100 400 20 10 11942 36 q 780 100 20 850 20 — 850 28 200 50 100 400 20 10 11942

2. Steel Structure

With respect to each sample, the ferrite fraction, the total fraction of tempered martensite and tempered bainite (mentioned as “tempered M/B” in Table 3), the amount of retained (amount of retained γ), the half-width of the concentration distribution in the carbon-concentrated region, and the scattering intensity at the q value of 1 nm⁻¹ in X-ray small angle scattering were determined by the above-mentioned methods. In the measurement of the amount of retained austenite, a two-dimensional microfocused X-ray diffraction apparatus (RINT-RAPID II) manufactured by Rigaku Corporation was used. The results are shown in Table 3.

In this Example, the steel structure (balance structure) other than the steel structure shown in Table 3 is untempered martensite in samples excluding sample No. 9, or untempered bainite in sample No. 9.

TABLE 3 Steel structure Half-width of Scattering Amount Average concentration distribution intensity Tempered of size of of Mn in carbon- at q value Steel Ferrite M/B retained γ MA concentrated region of 1 nm⁻¹ No. No. % % % μm % by mass cm⁻¹ 1 a 0 72 17.3 0.67 *0.23 *2.42 2 a 0 72 17.7 0.90 *0.17 0.74 3 a 0 71 17.1 0.67 *0.22 *2.34 4 a *30.7 *40 16.2 0.80 *0.24 0.72 5 a *30.5 *36 16.8 0.87 *0.24 0.74 6 a *13.7 61 18.0 0.70 0.71 0.68 7 a 0 *53 17.5 *2.07 0.86 0.71 8 a 0 72 17.1 0.60 *0.16 0.71 9 a 0 *0 16.4 *2.07 *0.22 0.81 10 a 0 80 *4.8 0.53 0.74 0.66 11 a 0 71 17.2 0.60 0.80 *1.30 12 a 0 70 16.4 0.67 0.89 0.97 13 a 0 70 18.1 0.73 0.80 0.89 14 a 0 71 16.5 0.67 0.76 0.72 15 a 0 72 17.7 0.73 0.77 0.73 16 a *21 *49 16.3 0.83 0.73 *2.40 17 a 0 72 16.5 0.63 0.88 *1.38 18 a 0 70 17.9 0.80 0.84 0.74 19 a 0 74 *6.5 0.87 0.89 *1.56 20 a 0 69 *6.5 0.77 0.83 0.72 21 b 0 73 14.2 0.70 0.82 0.73 22 c 0 77 *5.6 0.73 0.72 0.73 23 d 0 60 *9.2 0.83 0.82 0.71 24 e *26 *50 13.5 0.77 0.72 0.71 25 f 0 *51 *6.9 *1.70 0.76 0.70 26 g 0 61 29.9 0.73 0.85 0.67 27 h 0 71 17.6 0.63 0.82 0.67 28 i 0 71 16.7 0.67 *0.16 0.74 29 j 0 70 16.9 0.87 0.79 0.71 30 k 0 73 14.3 0.60 0.68 0.71 31 l 0 73 13.5 0.87 0.70 0.67 32 m 0 73 12.3 0.67 0.75 0.71 33 n 0 70 16.5 0.60 0.88 0.74 34 o 0 69 17.5 0.90 0.67 0.68 35 p 0 70 16.8 0.57 0.76 0.75 36 q 0 70 15.8 0.87 0.75 0.69

3. Mechanical Properties

With respect to the resulting samples, YS, TS and EL were measured using a tensile tester, and YR and TS×EL were calculated. By the above-mentioned methods, the hole expansion ratio λ, the maximum forming height (forming height), the cross tensile strength (SW cross tension) of the spot welded portion, and the R5 tensile thickness reduction ratio were determined. The results are shown in Table 4.

TABLE 4 Properties Forming SW cross R5 tensile thickness YS TS EL TS × EL λ height tension reduction ratio No. Steel No. MPa MPa YR % MPa % % mm kN % 1 a 973 1177 0.83 19.7 23170 39 *18.4 13.7 *32 2 a 997 1204 0.83 19.4 23302 32 *17.2 15.4 57.1 3 a 1011 1226 0.83 18.9 23114 38 *18.4 14.8 *30.1 4 a 600 *940 *0.64 23.6 22184 42 *18.6 14.3 54.1 5 a 639 *964 *0.66 23.4 22558 27 *18.8 14.4 55.5 6 a 761 1112 *0.68 23.4 25988 25 21 14.8 54.4 7 a 1182 1485 0.80 16.6 24717 *12 21.4 14.2 57.6 8 a 1011 1220 0.83 18.7 22846 28 *17.1 14.1 56.9 9 a 719 *938 0.77 22.9 21480 *11 *18.4 13.7 56.2 10 a 873 1069 0.82 14.5 *15494 38 *15.5 14.2 58.4 11 a 988 1195 0.83 20.3 24247 33 21.4 15 *45.2 12 a 1036 1256 0.83 19 23884 29 21.5 15.5 54.3 13 a 1127 1365 0.83 18.4 25167 44 21.4 15 56 14 a 1135 1373 0.83 17.4 23894 23 21.1 15.1 54.8 15 a 1000 1206 0.83 20.5 24775 34 21.3 14.6 56.9 16 a 632 *949 *0.67 22.9 21732 41 20.9 14 *32 17 a 1066 1288 0.83 18.7 24039 29 21.4 14 *42.4 18 a 1032 1252 0.82 19.9 24948 30 21.4 13.5 57.5 19 a 810 994 0.82 17 *16947 26 *16.9 19 *38.4 20 a 1278 1590 0.80 10.7 *16983 28 *17 11.4 58.4 21 b 1023 1237 0.83 18 22326 31 21.3 17.9 53.6 22 c 984 1204 0.82 13.3 *16072 39 *16.1 22.5 58.2 23 d 1164 1468 0.79 12.9 *18918 45 *18.9 18.1 56.2 24 e 592 *921 *0.64 19.8 *18235 23 21.2 17.1 55.3 25 f 1222 1608 0.76 10.7 *17206 *14 *17.2 14.1 56.2 26 g 1180 1416 0.83 23.4 33197 45 21.4 *4.2 58.1 27 h 1084 1312 0.83 13.4 *17552 23 *17.6 13.6 57.7 28 i 1046 1271 0.82 17.8 22589 24 *16.9 14.9 55.5 29 j 1056 1279 0.83 19 24295 39 21.3 13.2 58.5 30 k 1023 1243 0.82 17.6 21874 40 21 17.7 54.5 31 l 1054 1279 0.82 16.7 21357 45 21.1 18.1 59.1 32 m 1048 1280 0.82 16.3 20839 36 21.1 17.6 56.9 33 n 1093 1323 0.83 18.1 23903 42 21.5 15.6 57 34 o 1097 1331 0.82 18.3 24316 34 20.8 14.6 54.1 35 p 1058 1283 0.83 18.6 23853 44 21.1 14.6 55.5 36 q 1084 1324 0.82 17.6 23333 35 21.2 15.3 58.7

The results of Table 4 will be considered. Samples Nos. 13, 15, 18, 21 and 28 to 36 are Examples that satisfy all requirements (composition, manufacturing conditions and steel structure) defined in the embodiment of the present invention. All of these samples achieve the tensile strength (TS) of 980 MPa or higher, the yield ratio (YR) of 0.75 or more, TS×EL of 20,000 MPa % or higher, the hole expansion ratio (λ) of 20% or more, the maximum forming height of 16 mm or more, the SW cross tension of 6 kN or higher, and the R5 tensile thickness reduction ratio (RA) of 50% or more.

To the contrary, in sample No. 1, since holding was performed only at a temperature of an Ac₃ point or higher corresponding to a temperature T₂ without performing austenitizing in two stages of austenitizing at a temperature T₁ and austenitizing at the temperature T₂, the sample exhibits small half-width of the concentration distribution of Mn in the carbon-concentrated region and low maximum forming height. Since the holding time [7] was as long as 300 seconds, carbide (cementite) was precipitated. Because of large scattering intensity of X-ray small angle scattering, it can be said that cementite having a size of about 1 nm has a large volume fraction. As a result, collision resistance properties (thickness reduction ratio) were degraded.

Sample No. 2 exhibited small half-width of the concentration distribution of Mn in the carbon-concentrated region and low maximum forming height because of low holding temperature T₁.

Sample No. 3 exhibited small half-width of the concentration distribution of Mn in the carbon-concentrated region and low maximum forming height because of high holding temperature T₁. Since the holding time [7] was as long as 300 seconds, carbide (cementite) was precipitated. Because of large scattering intensity of X-ray small angle scattering, it can be said that cementite having a size of about 1 nm has a large volume fraction. As a result, collision resistance properties (thickness reduction ratio) were degraded.

Regarding samples Nos. 4 and 5, after heating to the heating temperature T₁ and holding, the temperature that is the same as T₁ was selected as the heating temperature T₂, thus failing to austenitize at sufficiently high temperature. Therefore, samples exhibited a large amount of ferrite, low total fraction of tempered martensite and tempered bainite, and small half-width of the concentration distribution of Mn in the carbon-concentrated region. As a result, samples exhibited low tensile strength, low yield ratio and low maximum forming height.

Sample No. 6 exhibited large amount of ferrite because of low heating temperature T₂, leading to low yield ratio.

Sample No. 7 exhibited low total fraction of tempered martensite and tempered bainite and large average size of MA because of high cooling stopping temperature T₃, thus decreasing the hole expansion ratio.

Sample No. 8 exhibited small half-width of the concentration distribution of Mn in the carbon-concentrated region because of short holding time at the heating temperature T₁, thus decreasing the maximum forming height.

Regarding sample No. 9, the holding time at the heating temperature T₂ was long and the cooling stopping temperature T₃ is high. Therefore, samples exhibited the total fraction of tempered martensite and tempered bainite of 0%, large average size of MA and small half-width of the concentration distribution of Mn, thus decreasing the tensile strength, the hole expansion ratio and the maximum forming height. Because of holding at the same temperature for 300 seconds (holding time [9]), the formation of carbide was suppressed, thus decreasing the hole expansion ratio λ.

Sample No. 10 exhibited small amount of retained austenite because of low cooling stopping temperature T₃, thus decreasing low value of TS×EL and maximum forming height.

Sample No. 11 precipitated carbide (cementite) since the reheating rate [8] was as low as 30° C./sec. Because of large scattering intensity of X-ray small angle scattering, it can be said that cementite having a size of about 1 nm has a large volume fraction. As a result, collision resistance properties (thickness reduction ratio) were degraded.

Sample No. 16 exhibited large amount of ferrite and low total fraction of tempered martensite and tempered bainite because of low rapid cooling starting temperature T₄, thus decreasing the tensile strength and the yield ratio. Since the holding time [9] was as long as 300 seconds, carbide (cementite) was precipitated. Because of large scattering intensity of X-ray small angle scattering, it can be said that cementite having a size of about 1 nm has a large volume fraction. As a result, collision resistance properties (thickness reduction ratio) were degraded.

Sample No. 17 precipitated carbide (cementite) since the reheating rate [8] was as low as 15° C./sec. Because of large scattering intensity of X-ray small angle scattering, it can be said that cementite having a size of about 1 nm has a large volume fraction. As a result, collision resistance properties (thickness reduction ratio) were degraded.

Sample No. 19 exhibited the parameter increased to 14,604 because of high reheating temperature T₅, thus decreasing the amount of retained austenite. As a result, the value of TS×EL and maximum forming height decreased. Because of large scattering intensity of X-ray small angle scattering, it can be said that cementite having a size of about 1 nm has a large volume fraction. As a result, collision resistance properties (thickness reduction ratio) were degraded.

Sample No. 20 exhibited the parameter decreased to 9,280 because of low reheating temperature T₅, thus decreasing the amount of retained austenite. As a result, the value of TS×EL and maximum forming height decreased.

Sample No. 22 exhibited small C amount, insufficient amount of retained austenite, thus decreasing the value of TS×EL and maximum forming height.

Sample No. 23 exhibited large Mn amount and insufficient amount of retained austenite, thus decreasing the value of TS×EL and maximum forming height.

Sample No. 24 exhibited small Mn amount, large amount of ferrite, and insufficient total amount of tempered martensite and tempered bainite. As a result, the tensile strength and the yield ratio decreased.

Sample No. 25 exhibited small amount of Si+Al, insufficient total amount of tempered martensite and tempered bainite, small amount of retained austenite, and large average size of MA. As a result, the value of TS×EL, the hole expansion ratio and the maximum forming height decreased.

Sample No. 26 exhibited large C amount, thus decreasing the SW cross tensile strength.

Sample No. 27 exhibited large amount of Si+Al, thus decreasing the value of TS×EL and the maximum forming height.

Sample No. 28 exhibited small half-width of the concentration distribution of Mn in the carbon-concentrated region and low maximum forming height because of high holding temperature T₁.

4. Conclusion

In this way, it could be confirmed that, regarding the steel sheet that satisfies the composition and the steel structure defined in the embodiment of the present invention, all of tensile strength (TS), yield ratio (YR), the product (TS×EL) of (TS) and total elongation (EL), hole expansion ratio (λ), thickness reduction ratio (RA) of the fracture portion during a tensile test, maximum forming height and cross tensile strength of the spot welded portion are at high level.

It could be also confirmed that the manufacturing method according to the embodiment of the present invention enables the production of the steel sheet that satisfies the composition and the steel structure defined in the embodiment of the present invention.

This application claims priority based on Japanese Patent Application 2016-153110 filed on Aug. 3, 2016, the disclosure of which is incorporated by reference herein. 

1. A high-strength sheet comprising Fe and: C: 0.15% by mass to 0.35% by mass, a total of Si and Al: 0.5% by mass to 3.0% by mass, Mn: 1.0% by mass to 4.0% by mass, P: 0.05% by mass or less, and S: 0.01% by mass or less, wherein the high-strength sheet comprises a steel structure wherein: a ferrite fraction is 5% or less, a total fraction of tempered martensite and tempered bainite is 60% or more, an amount of retained austenite is 10% or more, a martensite-austenite constituent has an average size of 1.0 μm or less, a half-width of a concentration distribution of Mn in a carbon-concentrated region that is equal to the amount of retained austenite is 0.3% by mass or more, and a scattering intensity at a q value of 1 nm ⁻¹ in X-ray small angle scattering is 1.0 cm⁻¹ or less.
 2. The high-strength sheet of claim 1, comprising 0.30% by mass or less of C.
 3. The high-strength sheet of claim 1, comprising less than 0.10% by mass of Al.
 4. A method for manufacturing a high-strength sheet, the method comprising: preparing a rolled material comprising Fe and: C: 0.15% by mass to 0.35% by mass, a total of Si and Al: 0.5% by mass to 3.0% by mass, Mn: 1.0% by mass to 4.0% by mass, P: 0.05% by mass or less, and S: 0.01% by mass or less, holding the rolled material, which has an Ac₁ point and an Ac₃ point, at a temperature between the Ac₁ point and 0.2×the Ac₁ point+0.8×the Ac₃ point for 5 seconds or more, followed by heating to a temperature of the Ac₃ point or higher and further holding for 5 to 600 seconds, thereby austenitizing the rolled material, after the austenitizing, cooling the material from a temperature of 650° C. or higher to a cooling stopping temperature between 100° C. or higher and lower than 300° C. at an average cooling rate of 10° C./sec or more; heating the material from the cooling stopping temperature to a reheating temperature T in a range of 300° C. to 500° C. at an average heating rate of 30° C./sec or more; holding at the reheating temperature T for a holding time of 1 to 150 seconds so as to satisfy a tempering parameter P of 10,000 to 14,500; and after the holding at the reheating temperature T, cooling from the reheating temperature T to 200° C. at an average cooling rate of 10° C./sec or more, wherein P=T×(20+log(t/3600))   (1) where T: reheating temperature (K) and t: holding time (seconds).
 5. The method of claim 4, wherein cooling to the cooling stopping temperature comprises: cooling to a rapid cooling starting temperature that is a temperature of 650° C. or higher at an average cooling rate of 0.1° C./sec or more and less than 10° C./sec; and cooling from the rapid cooling starting temperature to the cooling stopping temperature at an average cooling rate of 10° C./sec or more.
 6. The method of claim 4, wherein the tempering parameter P is in a range of from 11,000 to 14,000 and the holding time t is in a range of from 1 to 150 seconds. 